Ultramicroscopy 37 (1991) 169-179 169 North-Holland

Characterization of thin films, interfaces and surfaces by high-resolution electron microscopy

David J. Smith ], Z.G. Li, Ping Lu, M.R. McCartney and S.-C.Y. Tsen Center for Solid State Science, Arizona State University, Tempe, A Z 85287, USA

Received 7 January 1991

Recent applications of high-resolution electron microscopy in our laboratory to the characterization of thin films, interfaces and surfaces are described. The information typically available using this technique about thin films is illustrated by studies of multiple quan tum wells, metallic superlattices, magnetic films and X-ray optical elements. The problems associated with imaging interfaces are discussed by reference to silicon/silicide interfaces and grain boundaries in metals. Finally, recent observations of semiconductor surface reconstructions and beam-induced reactions at oxide surfaces are briefly summarized.

1. Introduction

The latest generation of high-resolution elec- tron microscopes (HREMs) have structural resolu- tion limits on the scale of atomic dimensions, making it possible to observe directly the atomic structure of crystal defects and interfaces [1,2]. In this short review, we describe some of our recent studies of thin films, interfaces and surfaces, and thereby illustrate the potentialities of the H R E M for determining unknown structures and solving important structural problems. The interested reader is referred elsewhere for further informa- tion about the theory and practice of high-resolu- tion electron microscopy [3,4].

2. Interface and surface imaging

Various types of interfaces between two dissim- ilar materials are illustrated schematically in fig. 1. These include, but obviously are not limited to: (a) abrupt interface between crystalline materials of (effectively) identical lattice parameter;

~ Also at Department of Physics.

(b) abrupt interface between crystalline materials of different lattice parameter; (c) abrupt interface between crystal l ine/amor- phous materials; (d) diffuse interface. (Of course, the interface between a surface and the adjacent atmosphere or vacuum should be consid- ered as a special type of interface.) Examples of these various types are given in the applications which follow. The features of practical interest

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fuse.

0304-3991/91/$03.50 © 1991 - Elsevier Science Publishers B.V. (North-Holland)

17(/ D.J. Smith et al. / ('haractertzat*on of thin film,s; ~nter]aces and xurfaces

include the abruptness of the interface, whether it is smooth, flat or facetted, the extent of inter- penetration of materials, and whether there are any orientational relationships (epitaxy) between the two materials. The existence of a further amorphous or crystalline phase is also possible in some interracial systems, and the presence of a surface oxide layer or the occurrence of a surface (or interface) reconstruction often needs to be considered. In compound materials, the identity of the terminating atomic species can have a marked influence on the surface or interfacial properties, and the nature of any defects required to accom- modate interfacial mismatch may be important. Finally, any structural differences introduced by the particular preparation method need to be analyzed and understood.

The major requirement for characterization of planar or layered structures by high-resolution electron microscopy is for the plane(s) of interest to be aligned with the incident beam direction. This is not a problem for multilayer materials when well defined surfaces of crystalline materials, such as GaAs or Si, provide the substrate for deposition, since these materials can also be utilized for orientation of the interface(s). Low-in- dex zones such as [100] and [110] are typically used, but [111] and [112] can provide further use- ful projections when three-dimensional views are required in order to specify the interface struc- tures fully. Observation in this cross-sectional mode, rather than with the electron beam normal to the plane(s) or interface(s), enables the micro- structure of the thin films or multilayers to be characterized as a function of lateral position as well as of distance from the substrate (or previous layer). Nevertheless, it must be appreciated that. in common with other transmission imaging modes, the features of the final HREM image depend upon the projection of the crystal struc- ture, normally presumed to be periodic in the beam direction. Sidesteps, other types of disorder, or even regular stacking of two or more species, in the forward direction cannot be distinguished, and very little useful information can be derived about the general curved or inclined grain boundary.

For high-resolution imaging rather than dif- fraction studies of surfaces, the microscope oper-

ating mode of choice is the technique of surface profile imaging [5]. The geometrical relationship between the specimen and the electron beam of the microscope is effectively the same as outlined above for interfaces, so that it is also possible to image surface features in profile with atomic reso- lution. Moreover, with the assistance of an online image pickup and viewing system, a variety of dynamic surface activities can be recorded in real time without significant loss of performance, and later analyzed in detail. Several applications of the method will be described below.

3. Applications

3.1. Thin films

Ultrathin multilayer films with periodic varia- tions in composition on the nanometer scale have unique physical properties which are of great fundamental and technological interest. Their suc- cessful utilization in optical, magnetic and mag- neto-optic devices depends critically, however, on the quality of the films, in particular the layer uniformity (composition and thickness) and the interface roughness. X-ray diffraction, Rutherford scattering and Auger profile analysis are bulk characterization techniques commonly used in multilayer studies, but they provide structural in- formation averaged over the entire bulk material. High-resolution electron microscopy can provide additional complementary information about the local microstructure as well as a more immediate and direct view of the overall structure. We con- sider here several examples of multilayer films of metals and semiconductors, as grown by electron- beam evaporation, sputtering and molecular-beam epitaxy (MBE) methods: the schematic of fig. 1 provides a useful basis for consideration. Recent reviews of MBE [6], metallic multilayers [7], and semiconductor superlattices [8] should be con- sulted for further background information and extensive lists of references.

3.1.1. Multtple quantum wells Ultrathin multilayers of I I I - V compound semi-

conductors such as GaAs/AIAs, possibly inter-

D.J. Smith et al. / Characterization of thin films, interfaces and surfaces 171

mixed with InP, have interesting device applica- tions. These so-called multiple quantum well (MQW) structures, generally grown by MBE, rep- resent a substantial challenge to the electron mi- croscopist because of their well-matched unit cells: they closely fit the schematic of fig. la. For exam- ple, under the imaging conditions traditionally required for high-resolution structure imaging, namely thin crystals and optimum defocus, it is extremely difficult to pinpoint the location of the interface, and hence virtually impossible to de- termine the interface sharpness. In order to high- light the compositional variations it is necessary to find thickness and defocus combinations which accentuate differences in the images from the two constituent materials.

Initial studies of MQWs utilized the (110) pro- jection, basically because of the available micro- scope resolution, but more attention has lately been given to imaging in the (100) projection because recent gains in microscope performance have made it possible to obtain structure images. Both projections feature {200} beams which are highly sensitive to chemical variations because

their amplitude depends upon differences in struc- ture factor. Since f~a --fAs is relatively small, the GaAs (200} beams remain weak up to substantial crystal thicknesses unlike, for example, those of AlAs. The (100) projection has 4 of these chem- ically sensitive (002} beams, compared with (110) which has only 2, and, unlike (110), there is no chance of double diffraction in (100) giving rise to misleading intensity [9]. Fig. 2 shows an exam- ple of GaAs/AIAs MQW imaged in the [001] projection, and the successive layers of AlAs (light) and GaAs (dark) are clearly visible. Image simula- tions for GaAs /Ga l_xAl~As indicate that inter- faces should still be distinct in (100) for values of x as small as 0.2 whereas interfaces are unlikely to be visible in (110) even for x less than about 0.3 [10]. Image contrast can be enhanced in (110) (or (100)) by deliberately restricting the image forma- tion process to only five beams with an objective aperture, despite some deterioration in available resolution [11,12]. Finally, as a guide to selecting the optimum imaging conditions, it can be helpful to compute the so-called generalized transfer func- tion which takes into account the relative phases of the (000) and {002} beams as well as the microscope transfer characteristics [13].

Fig. 2. Cross-sectional view of multiple quantum well of GaAs/A1As imaged in the [001] projection showing layers of

GaAs (dark) and AlAs 0ight).

3.1.2. Metal l ic superlattices

Relative to multilayers based on semiconduct- ing materials, metallic superlattices have received comparatively little attention, mainly because the former have had immediate practical applications and the latter are far more difficult to grow due to their lack of covalent bonding [7]. These most commonly and closely fit the example drawn in fig. lb. It is also unclear to what extent some sort of chemical mixing (i.e., interdiffusion) also occurs in some systems. Nevertheless, metallic multi- layers with a high degree of structural coherence, primarily grown by MBE on GaAs or Si sub- strates, are becoming more commonplace [6].

Multilayers based on transition-metal ferro- magnets, such as Fe and Co, display some intrigu- ing magnetic properties, atypical of bulk behavior, which have potentially important implications for recording media. For example, muitilayers consist- ing of C o / P d and C o / P t exhibit perpendicular magnetic anisotropy provided that the Co layer

172 D.J. Smith et aL / ('haracterication of thtn films, znter]ace,v and ~u(l~cev

Fig. 3. High-resolution electron micrograph showing buffer layers of Co (0.84 nm)/Ag (20 nm) grown by MBE on

GaAs(100) substratc.

thickness is kept below a critical value, typically on the order of 1 nm, which is apparent ly depen- dent upon the orientation of the GaAs substrate [14]. Unwanted reactions between the substrate and the multilayer constituents represent a prob- lem which is being addressed by using inter- mediate buffer films. Fig. 3 shows an example of a GaAs(100) substrate upon which a buffer layer of Co (nominal 0.84 n m ) / A g (nominal 20 nm) has been grown. Three to five layers of Co with a metastable bcc structure, as deduced by lattice spacing measurements, are clearly visible im- mediately adjacent to the substrate, and disloca- tions (arrowed) can be identified in the Ag layer by viewing the image at an inclined angle. These presumably accommoda te the strain which arises from the mismatch between the lattices of the different materials (about 2%). Magnetic measure- ments show that the multilayer properties are strongly related to the deposit ion conditions, and further studies are in progress to optimize the growth condit ions in relation to the microstructure and magnetic behavior.

ness and composit ion, in particular how these are influenced by the substrate temperature and the rate of deposition. In the case of Co/Cr multi- layers, which again are candidates for perpendicu- lar magnetic recording, a c o m m o n observation is the presence of a co lumnar structure which per- sists through a considerable number of bilayer periods [15]. However, as shown in fig. 4, pro- nounced layer contrast and coherent growth be- tween layers, visible as continui ty of lattice planes, may still be retained. Depending on the materials~ these films should ideally be as sketched in figs. la and 1 b.

3.1.4. X-r~ O' optical elements Recently, increased attention has been given to

ultrathin multilayers suitable for utilization in X- ray optical applications. These generally consist of alternating layers of high density, high-atomic-

3.1.3. Eoaporated magnetic films Electron-beam evaporat ion is another deposi-

tion technique commonly used for growing mag- netic thin films [7]. Morphological features of in- terest which can be studied by high-resolution electron microscopy include the intrinsic structure of the individual layers, and the interfacial thick-

Fig. 4. Metallic multilayer specimen grown by e-beam evapora- tion with nominal Co/Cr layer thicknesses of 2.8 nm/2.9 nm

respectively.

D.J. Smith et al. / Characterization of thin films, interfaces and surfaces 173

number elements (such as W, Mo and Rh) and low-density, low-atomic-number elements (such as C or Si), with bilayer periods which are ap- propriate for Bragg reflection at the particular wavelengths of interest. Fig. lc is the best repre- sentation of the ideal case, but fig. ld is more typical. In practice, reflectivity coefficients are typically less than half the expected theoretical values because of various factors associated with growth of the multilayers. For example, the two primary growth techniques for X-ray elements are sputtering and electron-beam evaporation. Sputtering is a comparatively energetic process: sputtered atoms have considerable energy when they strike the substrate, enabling them to migrate and thereby avoid three-dimensional island growth. However, interdiffusion between layers can then be a problem, particularly for the higher density metal into the lower-density material. Electron beam evaporation imparts much less en- ergy into the atoms being evaporated compared with sputtering, which is fine for limiting diffu- sion. However, it is not clear what effect the surface roughness of each successive layer has on the deposition of the subsequent layer.

Several high-resolution studies of bilayer sys- tems grown by either technique have recently been published (see, for example, refs. [16-19]). Ob- servations of W / C and Mo/S i systems grown by sputtering have confirmed the considerable dif- ferences in interface roughness which depend upon the relative atomic number of the two species. It is also interesting that, for a number of these bilayer systems, there appears to be a critical thickness for nucleation of metal crystallites. For example, W layers remain amorphous up to - 40 ~, thickness [16]. For W / C systems the growth of WC crystal- lites has been observed [17], and the collapse a n d / o r separation of the W layers due to C sputtering by the electron beam during observa- tion has also been reported [17]. Recent studies of M o / S i bilayer materials grown by electron-beam evaporation have concentrated upon the depen- dence of the structural properties on the deposi- tion rate and the substrate temperature [18], and it was shown that interpenetration of Mo into Si could be restricted by using comparatively large angles of evaporation (as measured with respect to

Fig. 5. Mo/Si multilayer evaporated at substrate temperature of ~ K with deposition angle of 42 ° relative to substrate

normal.

the substrate normal). Fig. 5 shows a crystal grown at a substrate temperature of 400 K with an incident deposition angle of 42 ° , as viewed in the direction normal to the incident evaporated beam, so that the columnar growth of tilted Mo crystal- lites is clearly visible. A comparison of carbon and boron carbide as layer spacing materials for multi- layers grown by sputtering has shown that the latter result in lower interface roughness [19]. Op- t imum growth conditions for other important sys- tems have not yet been established, and further experimentation, allied with H R E M observations, is still required before broad generalizations can be made.

3.2. Interfaces

The interface between two similar or dissimilar materials has a profound influence on the behav- ior of the composite material. A detailed knowl- edge of the atomic structure and chemical com- position in the vicinity of the interface should enable invaluable insight into its physical proper-

174 D.J. Smith et al. / ('haracterization of thin films, interfaces and sur]acex

ties to be gained. High-resolution electron mi- croscopy therefore has a valuable role to play in the characterization of interfaces. The require- ments for determination of atomic arrangements. assuming both materials are crystalline, are again simply stated. Both crystals should be aligned in low-index zone axes, the interface should be aligned so that it is edge-on to the incident beam, and it should be flat, i.e., there should be no sidesteps normal to the beam direction. These conditions can be met for many special types of planar faults including twins, Guinier-Preston zones and crystallographic shear defects. More- over, provided that the incident beam is well aligned with the optic axis of the objective lens, it may be possible to locate the relative positions of atomic columns at an interface to an accuracy approaching 0.1 ~ [20]. The opportunity for HREM to have an impact on interface studies is illustrated here with reference to a silicon/silicide boundary and a low-angle tilt grain boundary in AI.

3.2.1. Silicon/silicide interfaces Several metal silicides (e.g., NiSi 2, CoSi2~ ErSi 2

and PtSi2) are known to grow epitaxially on sili- con. Since the structure at the interface is likely to affect the associated Schottky barrier height (SBH), and hence influence any electronic applica- tions based on the interface, detailed characteriza- tion at the atomic level is likely to be highly worthwhile. The major question to be resolved is the local coordination of the metal atoms at the interface (possibly 5-fold, 7-fold or 8-fold) since

this determines the number of dangling bonds and ultimately the SBH. Moreover, there is also a possibility of two types of epitaxial growth, known as A-type (lattice planes aligned with those of silicon) or B-type (lattice planes in twinned orien- tation). Detailed analysis of A- and B-type struc- tures in NiSi 2 enabled structural models to be established [21], and it was later shown that there was a 0.2 eV difference in the SBH for the two orientations [22].

In the particular case of the CoSi2/Si (111) interface, both types of silicide orientation are known to occur [23], and it is again important to establish the Co coordination at the interface for each. Possible interface atomic configurations for 5-fold, 7-fold, and 8-fold coordination can be pro- posed for both A- and B-type interfaces [24]. The problem for the microscopist, in common with those experienced with MQWs, is to locate imag- ing conditions which enable differentiation be- tween these model structures to be made. Image simulations which incorporate the microscope parameters prove to be essential for this task. These confirm, for example, that the ~d [ l l l ] dis- crepancy in rigid body shift between the 7-fold and 5-fold (or 8-fold) models will be easily visible. Moreover, despite more subtle differences between the images for 5-fold and 8-fold models, suffi- ciently characteristic contrast features can be found which permit an unambiguous determina- tion of the interface structure [24].

A further significant aspect of silicon/silicide interfaces is the process whereby the actual silici- dation takes place. High-resolution observations

Fig. 6. Ledges with sequences of [111] interface steps at a CoSi2/Si(111) interface.

D.J. Smith et al. / Characterization of thin films, interfaces and surfaces 175

reveal a high density of [111] interfacial steps, and fig. 6 shows a stepped region of the CoS i2 /S i ( l l l ) interface which had been annealed at 400°C. Fur- ther in situ annealing experiments are required to investigate the details of the reaction mechanisms, but it appears that these [111] steps could provide a convenient means whereby nucleation and atomic intermixing occur at the interface.

3.2.2. Grain boundaries in metals A detailed knowledge of the core structure of

grain boundaries is necessary for understanding the behavior of individual interfaces although it may be of limited value for the general polycrys- talline material where the behavior is far more complicated. Lack of structural order in the beam direction imposes a restriction on the applicability of high-resolution electron microscopy to the gen- eral, arbitrarily oriented, grain boundary. Never- theless, recent improvements in microscope per- formance have made it possible to characterize atomic arrangements at (pseudo-)periodic grain boundaries in certain metals, at least in low-index

zones. Early studies included the characterization of interfacial dislocations at Z41 grain boundaries in molybdenum [25], and the analysis of coherent and incoherent (110) tilt boundaries in gold [26].

Several recent investigations have considered symmetrical and asymmetrical boundaries in aluminium [27-29]. The core structures of 60 ° and Lomer dislocations in small-angle asymmetric [110] tilt boundaries were established with the assistance of extensive image simulations [27], and the periodic recurrence of basic structural units was observed in a 90 ° symmetrical grain boundary [28]. Imaging in [100] projections is more demand- ing of the HREM, but fig. 7 shows an example of a symmetric tilt, low-angle (6 ° ) boundary in aluminium. The atomic arrangements of the regu- larly spaced array of dislocations along the inter- face are clearly visible, and analysis establishes that the associated Burgers vectors are ½1110], as expected from simple dislocation theory [29].

3.3. Surfaces

Fig. 7. A symmetric [100] tilt grain boundary in aluminium. Burgers vector of dislocation is b = -~ [110].

Profile imaging is a technique capable of pro- viding direct information about surfaces and surface reactions on an atomic scale, but the surface in profile should be thin ( < 100 ,& typi- cally) and free of contamination or reaction prod- ucts, and the H R E M must be operated correctly (i.e., proper adjustment of focus, image astigma- tism and beam alignment). These diverse require- ments have been satisfied for a range of semicon- ductors, oxides and metals [5]. However, operation under ultrahigh vacuum, with facilities for in situ surface treatment, is required before some image features, expecially those involving reactive surfaces, can be regarded as representative. We have also found, as shown below, that it can be particularly informative and worthwhile to under- take surface profile imaging while the specimen is held at elevated temperature. We briefly describe here our recent observations of surface reconstruc- tions for (100)CdTe and electron-beam-stimulated effects in SnO 2 and TiO 2. The interested reader is referred elsewhere for a recent review of surface profile imaging and an extensive list of references [5].

176 D.J. Smi th et a L / ( 'haracterieation of thin f i lms, interfaces and xur~ce.s

3.3.1. Semiconductor reconstructions Termina t i on of the bulk lat t ice of a semicon-

duc to r p roduces many " 'dangl ing" (unsat isf ied) bonds , which result in an uns table conf igura t ion that may cause subsequent rear rangements , which could be subs tant ia l , of the surface and near- surface a toms to take place. Because of their in- herent reactivity, s emiconduc to r surfaces are dif- ficult to keep clean par t i cu la r ly in the highly con- fined space of the object ive lens pole-pieces and spec imen stage where p u m p i n g speeds and vacuum levels are not so good. Nevertheless , some recent prof i le image studies of surface recons t ruc t ions have been repor ted [30 32]. In situ heat ing of Si up to its mel t ing po in t under ul t rahigh vacuum cond i t ions achieved surface cleanliness and a novel (1 x 1) d imer recons t ruc t ion of a (113~ surface was observed [30]. We ob ta ined clean surfaces of CdTe by focussing an intense e lect ron beam onto a small area, and recons t ruc t ions of (001), (110) and (111) surfaces resul ted [31]. However , in this case, the vacuum level was only - 10 ~ Tor r and uncer ta in ty thus remained over whether these re- cons t ruc t ions were representa t ive of par t ia l ly con- t amina ted , ra ther than clean, surfaces.

We have since repea ted our observa t ions of (001) CdTe surfaces under subs tan t ia l ly improved vacuum cond i t ions ( - 3 × 1 0 9 Torr) , using in situ annea l ing at - 500°C to ini t ia te sub l imat ion of mater ia l f rom edges of the bulk crystal [32]. Charac te r i za t ion of the surface recons t ruc t ions was then poss ib le once the crystal had been cooled to t empera tu res below abou t 300°C. It was found that the (001) surface had a (2 × 1) recons t ruc t ion at t empera tu res up to abou t 200°C which trans- formed re~ersibl, v into a (3 × 1) recons t ruc t ion over the a p p r o x i m a t e t empera tu re range of 200°C < T < 300°C. Profile images which d isp lay these recons t ruc t ions are shown in figs. 8a and 8b. St ructura l models for the (2 × 1) and (3 × 1) re- cons t ruc t ions were ob t a ined direct ly on the basis of the exper imenta l images and then conf i rmed by image s imulat ions . The (2 × 1) recons t ruc t ion in- volves a 1 / 2 mono laye r of Cd vacancies and a very large inward con t rac t ion of the remain ing Cd surface atoms, which then d isp lace the second layer of Te atoms. The (3 × 1) recons t ruc t ion in- volves both the fo rmat ion of surface dimers and

a ~ ~ ~

~

Fig. 8. Profile images from (001) CdTe surface showing (a) (2x1). and (b) (3x l) reconstructions, recorded at tempera-

tures of about 140°( ̀ and 2400( ", respectivel,,..

the presence of vacancies at the surface. Every third a tomic -pa i r is missing a long the [1, 1, 0] direct ion, and the remain ing two a tom pairs at the surface form the surface dimer. Both of these recons t ruc t ions were not seen in the ear l ier work under poore r vacuum condi t ions .

A tomic mot ion on the CdTe surfaces could be observed with the image-p ickup sys tem a t tached to the microscope. Wi th the in situ hea t ing facili ty, the t empera tu re dependence of a tomic mot ion could be s tudied and it was found that the dy- namic behavior clear ly visible on the surfaces was face-dependent . In general , the (001) surface was much more active than the (110) and (111) surfaces at all t empera tures , and any s t ruc tura l charac ter i - za t ion for this surface was real ly only poss ib le for t empera tu res be low 300°C. The (111) surface layers were of ten removed layer -by- layer in a pro- cess involving a ledge mechan i sm in which the removal was in i t ia ted at the pos i t ion of surface steps.

3.3.2. Beam-induced reactions in oxides During ex tended H R E M observa t ions of oxide

surfaces , a r ange of e l e c t r o n - b e a m - i n d u c e d processes can take place [5]. These inc lude surface diffusion, as well as surface, and near-surface, reduct ion due to e lec t ron-s t imula ted desorp t ion

D.J. Smith et al. / Characterization of thin films, interfaces and surfaces 177

(ESD) of oxygen. In some maximally-valent tran- sition metal oxides, the epitaxial development of a lower (metallic) oxide is observed [33], but a variety of end-products can in general be anticipated de- pending on such factors as the initial surface cleanliness, the residual microscope vacuum and the degree of covalency/ionici ty of the original oxide [34,35]. We describe examples here where the local current density of the incident electron beam and the temperature of the irradiated sam- ple produce further unexpected results.

Using the focussed probe of an analytical elec- tron microscope equipped with a field emission gun (FEG), current densities on the order of 103- 10 4 A / c m 2 can be achieved (as compared with 5-50 A / c m 2 in normal H R E M operating mode). Electron irradiation under such extreme condi- tions usually results in amorphization, reduction and sputtering from the irradiated region, with fringes due to the base metal occasionally ob- served within the irradiated area [35]. Under con- ventional H R E M imaging conditions, surfaces of

Fig. 10. Image of facetted holes in rutile single crystal recorded at a temperature of - 5 8 0 ° C following electron irradiation

( - 80 A / c m 2, 300 keV).

Fig. 9. High-resolution image of [001] SnO 2 showing tin crys- tallite which developed during intense electron irradiation ( -

10 4 A / c m 2, 100 keV).

SnO 2 did not develop crystalline regions of re- duced oxide, as had been observed for its more ionic isomorph TiO2, although substantial elec- tron-beam-induced facetting was observed. Under intense irradiation, however, regions of SnO 2 were reduced to perfect single crystals of metallic tin. Fig. 9 shows a high-resolution image of such a particle taken after the sample had been trans- ferred to a microscope of better resolution after irradiation in a FEG-equipped microscope [36].

Under (nominal) room-temperature operating conditions, it is clear that energy transfer from the incident electron beam is responsible for a sub- stantial portion of the physical and chemical processes observed to take place at oxide surfaces. It is therefore of considerable interest to observe changes in behavior which result from concurrent specimen annealing. Accordingly, we have re- cently investigated electron-beam-induced damage and annealing processes in rutile crystals, utilizing a 300 keV H R E M equipped with a heating holder

178 l).J. Smt th et aL / ('haractertz_atton ql thtn.filtr~, tntetjacex and xuU~wes

which had been m o d i f i e d for u l t rah igh v a c u u m

[37]. I r r a d i a t i o n at m o d e r a t e t e m p e r a t u r e s (200

4 0 0 ° C ) inc reased the ra te of su r face d a m a g e and

resu l ted in m u c h m o r e p r o n o u n c e d spu t t e r pi ts in

b o t h fu t i le and an i n t e r m e d i a t e TiO2-1I phase .

H o w e v e r , as s h o w n in fig. 10, i r r ad i a t i on at t em-

p e r a t u r e s in excess of 5 0 0 ° C resu l ted in the f o r m a -

t ion of well f ace t t ed holes in the ru t i le but,

surpr i s ing ly , no e v i d e n c e for any r educed phases

was obse rved .

In a n a r r o w sense, e l e c t r o n - b e a m - i n d u c e d reac-

t ions such as those d e s c r i b e d a b o v e can be cons id -

e red as e x t r a n e o u s : the phys ica l and chemica l

c h a n g e s to the ma te r i a l u n d e r o b s e r v a t i o n shou ld

be kep t to a m i n i m u m . Yet nove l ins ights in to

E S D m e c h a n i s m s for t r ans i t ion me ta l ox ides have

been o b t a i n e d us ing H R E M [33,34] and , in a

b r o a d e r sense, it a p p e a r s that in situ o b s e r v a t i o n s

cou ld help to be t t e r u n d e r s t a n d reac t ion m e c h a -

n i sms that take place , for example , d u r i n g nano -

l i t hography . W e be l ieve that H R E M has an im-

p o r t a n t role to p l ay in this i n t e re s t ing area of

su r face reac t ions .

4. C o n c l u s i o n

H i g h - r e s o l u t i o n e l ec t ron m i c r o s c o p y c o n t i n u e s

to r ep re sen t a power fu l t e c h n i q u e for e luc ida t i ng

the a t o m i c s t ruc tu re of th in fi lms, in te r faces and

surfaces . W i t h cu r r en t r e so lu t ion l imits of 0.15

0.20 nm, the t e c h n i q u e is be ing successful ly ap-

p l ied to a who le hos t of mater ia l s , r ang ing f rom

ce ramics , me ta l s a n d s e m i c o n d u c t o r s to mine ra l s

and o the r i n o r g a n i c oxides . Q u a n t i t a t i v e c o m p a r i -

sons wi th a tomis t i c mode l s , d e p e n d e n t u p o n

images r e c o r d e d in several p ro jec t ions , a re still

be ing in i t i a ted bu t i m p r o v e d reso lu t ion is needed

to ut i l ize images r e c o r d e d in h i g h e r - o r d e r zone

axes. In this regard , it is not o b v i o u s that the t r a d i t i o n a l r ou t e of r e so lu t i on i m p r o v e m e n t ,

n a m e l y by i nc rea s ing the acce l e r a t i ng vol tage , will

be benef ic ia l s ince the l i ke l ihood of a t o m i c dis-

p l a c e m e n t s , pa r t i cu l a r ly in the v ic in i ty of in te r -

faces and surfaces , is g rea t ly increased . I m a g e

r e s to ra t ion t echn iques , p e r h a p s based u p o n ho-

lography , m a y p r o v e to be m o r e e f fec t ive in the

long term. T h e p r o v i s i o n of in situ a n n e a l i n g

facil i t ies, as r epo r t ed here and e l sewhere [38], as

well as an u l t r ah igh v a c u u m e n v i r o n m e n t m the s p e c i m e n region, can also e n a b l e new ins ights in to

i m p o r t a n t phys ica l and c h e m i c a l p rocesses at

sur faces to be o b t a i n e d . F u r t h e r i n s t r u m e n t a l de-

v e l o p m e n t s are, however , still r e q u i r e d be fo re these

s tudies can be execu ted on a rou t ine , e v e r y d a y

basis.

Acknowledgements

This w o r k has r ece ived s u p p o r t f r o m the I B M

Sha red U n i v e r s i t y R e s e a r c h ( S U R ) P r o g r a m

and N a t i o n a l Sc ience F o u n d a t i o n G r a n t D M R -

8514583. W e are p leased to a c k n o w l e d g e m a n y

f r iends and co l l eagues for the i r o n g o i n g c o l l a b o r a -

t ion on s o m e of the p ro jec t s d e s c r i b e d here.

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